Materials Processing Technology ELSEVIER Joumal of Materials Processing Technology 190(2007)358-362 www.elsevier.com/locate/jmatprotec Rapid fabrication of c/C/Sic composite by Plp of HMds Jiehua Zhong Shengru Qiao a, Guofeng Lu, Yuebing Zhang a Ultra-high-temperature structural composites laboratory, Northwestern Polytechnical University, Xi'an 710072, PR China National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, PR China Received 9 July 2006: received in revised form 10 December 2006: accepted 3 February 2007 Abstract The C/C/SiC composite was fabricated within several days by the method of precursor impregnation and pyrolysis(PlP)using hexamethyldis- ilazane(HMDS)as the precursor. The carbon fiber plates, woven and punctured with two-dimensional orthogonal continuous carbon bundle and short carbon fiber, were used as the reinforced preforms. The characters of the C/C/SiC composite were analyzed using XRD, EDS and SEM, and three-point-bending test. The results indicated that pyrolyzed substance of the precursor contained excessive carbon, followed by silicon, and small amount of nitrogen. There were micro-cracks on the massive matrix. The phase composition was difficult to distinguish, primarily considered as the Si-C-N composite. The structure of C/C/SiC was dense and homogeneous with some pores whose sizes were less than 5 um between fibers and less than 100 um between the carbon fiber bundles, respectively Fiber pull-out was observed on the bending fracture surface. The stress-strain curve of both at room temperature and 1300 C appeared ascendin zigzag Flexural strength was 150 MPa at 1300C, higher than 121 MPa at room temperature o 2007 Published by Elsevier B.V. Keywords: PIP: HMDS; C/C/SiC; Flexural strength 1. Introduction Liquid precursor infiltration and pyrolysis(PIP) route to fab- ricate ceramic matrix composites is being actively studied for Continuous carbon fiber reinforced ceramic matrix com- its low processing temperature, short manufacturing time and posites(CFCCs) have many potential applications in high- simple demand for equipment [5,6]. Common precursors for temperature structural components because of their high strength fabricating C/SiC composites are such high molecular weight at elevated temperature, low density, superior toughness and polymers as polycarbosilane(PCs)and polymethlsilane(PMs brasive properties. Among CFCCs, carbon fiber reinforced However, their prices are high because they are not commer- silicon carbide composite(C/SiC) have been extensively investi- cial production at present. So, it is essential to look for a new gated and obtained certain practical application[1-3]. Currently, available precursor the main method for fabricating C/SiC composites in industry In this paper, HMDS, a cheap commercial organic substance, is the isothermal chemical vapor infiltration(CVD) technique. is selected. Feasibility of using it as precursor to fabricate C/Sic However, there was a limitation for its further development and composite is investigated and typical properties of as-prepare application because of its long price and long processing period C/C/SiC composite are analyzed [4]. In order to shorten processing period, several developments were done on CVI, including thermal-gradient chemical vapor 2. Material preparation analytical methods infiltration, forced-flow chemical vapor infiltration, pulse chem- ical vapor infiltration and so on. Those techniques need further In present work, the HMDS was used as precursor. The plates, woven and development although their processing time can be shortened. carbon fiber, were usedas the reinforced preform. The main characters of HMDS and the physical parameters of carbon fiber preform were given in Table 1.At preform was immersed in the container that was filled with molten low sponding author.Tel:+862988492084:fax:+862988492084 temperature coal pitch, and infiltrated with a pressure nailaddresses:zhongjiehua1981@mailnwpu.edu.cn, was taken out as the temperature dropped slightly and then put in a steel box. blao@nwpu.edu.cn(.Zhong) The preform, covered with a graphite powder layer with a thickness of 0.01 0924-0136/S-see front matter o 2007 Published by Elsevier B.V. doi: 10. 1016/j- jmatprotec 2007.02.008
Journal of Materials Processing Technology 190 (2007) 358–362 Rapid fabrication of C/C/SiC composite by PIP of HMDS Jiehua Zhong a, Shengru Qiao a,∗, Guofeng Lu a, Yuebing Zhang b, Wenbo Han b, Dechang Jia b a Ultra-high-temperature structural composites laboratory, Northwestern Polytechnical University, Xi’an 710072, PR China b National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, PR China Received 9 July 2006; received in revised form 10 December 2006; accepted 3 February 2007 Abstract The C/C/SiC composite was fabricated within several days by the method of precursor impregnation and pyrolysis (PIP) using hexamethyldisilazane (HMDS) as the precursor. The carbon fiber plates, woven and punctured with two-dimensional orthogonal continuous carbon bundle and short carbon fiber, were used as the reinforced preforms. The characters of the C/C/SiC composite were analyzed using XRD, EDS and SEM, and three-point-bending test. The results indicated that pyrolyzed substance of the precursor contained excessive carbon, followed by silicon, and a small amount of nitrogen. There were micro-cracks on the massive matrix. The phase composition was difficult to distinguish, primarily considered as the Si–C–N composite. The structure of C/C/SiC was dense and homogeneous with some pores whose sizes were less than 5 m between fibers and less than 100 m between the carbon fiber bundles, respectively. Fiber pull-out was observed on the bending fracture surface. The stress–strain curve of both at room temperature and 1300 ◦C appeared ascending zigzag. Flexural strength was 150 MPa at 1300 ◦C, higher than 121 MPa at room temperature. © 2007 Published by Elsevier B.V. Keywords: PIP; HMDS; C/C/SiC; Flexural strength 1. Introduction Continuous carbon fiber reinforced ceramic matrix composites (CFCCs) have many potential applications in hightemperature structural components because of their high strength at elevated temperature, low density, superior toughness and abrasive properties. Among CFCCs, carbon fiber reinforced silicon carbide composite (C/SiC) have been extensively investigated and obtained certain practical application [1–3]. Currently, the main method for fabricating C/SiC composites in industry is the isothermal chemical vapor infiltration (CVI) technique. However, there was a limitation for its further development and application because of its long price and long processing period [4]. In order to shorten processing period, several developments were done on CVI, including thermal-gradient chemical vapor infiltration, forced-flow chemical vapor infiltration, pulse chemical vapor infiltration and so on. Those techniques need further development although their processing time can be shortened. ∗ Corresponding author. Tel.: +86 29 88492084; fax: +86 29 88492084. E-mail addresses: zhongjiehua1981@mail.nwpu.edu.cn, blao@nwpu.edu.cn (J. Zhong). Liquid precursor infiltration and pyrolysis (PIP) route to fabricate ceramic matrix composites is being actively studied for its low processing temperature, short manufacturing time and simple demand for equipment [5,6]. Common precursors for fabricating C/SiC composites are such high molecular weight polymers as polycarbosilane (PCS) and polymethlsilane (PMS). However, their prices are high because they are not commercial production at present. So, it is essential to look for a new available precursor. In this paper, HMDS, a cheap commercial organic substance, is selected. Feasibility of using it as precursor to fabricate C/SiC composite is investigated and typical properties of as-prepared C/C/SiC composite are analyzed. 2. Material preparation analytical methods In present work, the HMDS was used as precursor. The plates, woven and punctured with two-dimensional orthogonal continuous carbon bundle and short carbon fiber, were used as the reinforced preform. The main characters of HMDS and the physical parameters of carbon fiber preform were given in Table 1. At first, the preform was immersed in the container that was filled with molten lowtemperature coal pitch, and infiltrated with a pressure of 1 MPa. The preform was taken out as the temperature dropped slightly and then put in a steel box. The preform, covered with a graphite powder layer with a thickness of 0.01 mm 0924-0136/$ – see front matter © 2007 Published by Elsevier B.V. doi:10.1016/j.jmatprotec.2007.02.008
J Zhong et al. Journal of Materials Processing Technology 190(2007)358-362 Table I The main characters of HMDS and physical parameters of carbon fiber preform Molecular structural formula (CH3 )3SiNHSI(CH3)3 61×41×11 Purity (%o) 990 Colorless and transparent liquid Bulk density (g/cm. 0.596 Preform Tolal Boiling point°C) Non-toxicity (g) specimen 6 Fig. 1. Schematic diagram for three-point-bending test(all dimensions were in mm). as innermost layer, followed by a fine quartz layer with a thickness of 0. 1 mm, a disordered fiber zone(zone B) in Fig. 4. The morphology of arse quartz sand layer with a thickness of 2-3 mm at outermost, was set into magnified zone B was showed in Figs. 5 and 6 From Fig. 6, it the pyrolysis batch and the pyrolysis was carried out at 950C with a pressure of was observed that the thickness of pyrocarbon layer was about 0.1 MPa, N2 used as protective atmosphere. Then a pyrocarbon layer was formed on the surface of carbon fiber of the preform. The preform with pyrocarbon layer 1.5 um as immersed into the container filled with HMDS, and pyrolysis was perf The zone C whose fibers were normal to fracture surface ubsequently. The pyrolysis processing was the same as that of pitch except that appeared the features of brittle failure. Smooth fracture and lit- pyrolysis temperature was 850C. The sample was densified by repeating four tle fiber pull-out were observed(Fig. 4). The morphology of cycles of infiltration and pyrolysis of HMDS. The overall procedure could be magnified zone C was showed in Fig. 5. It could be seen that Length direction of bend specimens was parallel to one of continuous carbon the thickness of pyrocarbon interface was about 0.5 um, and the xx 3.5 mm. Schematic diagram boundary between pyrocarbon layer and matrix was not clear, for three-point-bending test was illustrated in Fig. 1. The span for bending test which might suggest the strong interfacial bonding was 60mm. ASTM C1341-97 recommended that span-to-depth ratio was 16: 1 Fig. 7 was the magnification of zone A in Fig 4. There were However, when the span-to-depth ratio was larger than 8: 1, the test results would a lot of micro-cracks being approximately vertical to fiber and a be satisfactory [7] few cracks being parallel to fiber. These cracks partially formed 0.5 mm/min on a vacuum super-high-temperature testing machine at room tem. in the manufacturing process, and partially derived from damage perature and 1300C Flexural strength o was calculated from process during bending test. At the manufacturing temperature there were no cracks generating in the composite. However, when the composite was cooled to room temperature, owing to mismatch of the thermal expansion between fiber and matrix, where P is the maximum load before failure of the sample. L the length of span, b the width of the sample, h is the thickness of sample. The mor- phology of the sample was observed in the SEM. The chemical composition and the phase structure of matrix were examined with the EDS and XRD, 3. Results and discussion 3.1. The structure of the sample Figs. 2 and 3 showed that the macrostructure of as-prepare material was dense and homogeneous with some pores whose sizes were less than 5 um between fibers and less than 100 um between the carbon fiber bundles, respectively Because the plates used as the preform were woven and punctured with two-dimensional orthogonal continuous car- bon fiber and short carbon fiber. the laminated structure was observed on the fracture, each layer appeared different frac- ture characteristic (Fig. 4). There was fiber pull-out on the
J. Zhong et al. / Journal of Materials Processing Technology 190 (2007) 358–362 359 Table 1 The main characters of HMDS and physical parameters of carbon fiber preform HMDS Molecular structural formula (CH3)3SiNHSi(CH3)3 Preform Physicaldimension (mm) 361 × 41 × 11 Purity (%) ≥99.0 Characterization Colorless and transparent liquid Bulk density (g/cm3) 0.596 Flash point (◦C) 27 Tolal mass (g) 97 Boiling point (◦C) 126 Toxicity Non-toxicity Fig. 1. Schematic diagram for three-point-bending test (all dimensions were in mm). as innermost layer, followed by a fine quartz layer with a thickness of 0.1 mm, a coarse quartz sand layer with a thickness of 2–3 mm at outermost, was set into the pyrolysis batch and the pyrolysis was carried out at 950 ◦C with a pressure of 0.1 MPa, N2 used as protective atmosphere. Then a pyrocarbon layer was formed on the surface of carbon fiber of the preform. The preform with pyrocarbon layer was immersed into the container filled with HMDS, and pyrolysis was performed subsequently. The pyrolysis processing was the same as that of pitch except that pyrolysis temperature was 850 ◦C. The sample was densified by repeating four cycles of infiltration and pyrolysis of HMDS. The overall procedure could be finished within several days. Length direction of bend specimens was parallel to one of continuous carbon bundles. The specimen size was 70 mm × 5mm × 3.5 mm. Schematic diagram for three-point-bending test was illustrated in Fig. 1. The span for bending test was 60 mm. ASTM C1341-97 recommended that span-to-depth ratio was 16:1. However, when the span-to-depth ratio was larger than 8:1, the test results would be satisfactory [7]. Three-point-bending strength was measured with loading velocity of 0.5 mm/min on a vacuum super-high-temperature testing machine at room temperature and 1300 ◦C. Flexural strength σ was calculated from σ = 3PL 2bh2 where P is the maximum load before failure of the sample, L the length of span, b the width of the sample, h is the thickness of sample. The morphology of the sample was observed in the SEM. The chemical composition and the phase structure of matrix were examined with the EDS and XRD, respectively. 3. Results and discussion 3.1. The structure of the samples Figs. 2 and 3 showed that the macrostructure of as-prepared material was dense and homogeneous with some pores whose sizes were less than 5m between fibers and less than 100m between the carbon fiber bundles, respectively. Because the plates used as the preform were woven and punctured with two-dimensional orthogonal continuous carbon fiber and short carbon fiber, the laminated structure was observed on the fracture, each layer appeared different fracture characteristic (Fig. 4). There was fiber pull-out on the disordered fiber zone (zone B) in Fig. 4. The morphology of magnified zone B was showed in Figs. 5 and 6. From Fig. 6, it was observed that the thickness of pyrocarbon layer was about 1.5m. The zone C whose fibers were normal to fracture surface appeared the features of brittle failure. Smooth fracture and little fiber pull-out were observed (Fig. 4). The morphology of magnified zone C was showed in Fig. 5. It could be seen that the thickness of pyrocarbon interface was about 0.5m, and the boundary between pyrocarbon layer and matrix was not clear, which might suggest the strong interfacial bonding. Fig. 7 was the magnification of zone A in Fig. 4. There were a lot of micro-cracks being approximately vertical to fiber and a few cracks being parallel to fiber. These cracks partially formed in the manufacturing process, and partially derived from damage process during bending test. At the manufacturing temperature there were no cracks generating in the composite. However, when the composite was cooled to room temperature, owing to mismatch of the thermal expansion between fiber and matrix, Fig. 2. Cross section of the C/C/SiC composite.
J. Zhong et al. /Journal of Materials Processing Technology 190(2007)358-362 183sE ×1,seB1a Fig 3. The magnification of Zone C in Fig. 4 Fig. 6. The magnification of Zone B in Fig. 4 C B B A Fig 4. Fracture morphology of the C/C/SiC composite. Fig. 7. The magnification of Zone A in Fig. 4 thermal stress wound generate, which leaded to matrix cracking of composite. The micro-cracks could cause the deflection of 3.2. Analysis of phase structure and chemical composition the main crack when the material was subjected to load, which could improve the toughness of the composite. The XRD pattern showed in Fig 8 indicated that the matrix completely consisted of the amorphous substances, which was similar to the results in Refs. [8, 9]. Because the pyrolysis tem- which resulted in the non-crystalline structure of the matrix. The EDS pattern which was obtained from the position marked in Fig. 9 and the analysis result were given in Table 2, respectively It showed that the chemical compositions of matrix were mainly Si, C N and c seemed excessive according to the stoichiometric Elemental compositions of matrix of the as-prepared composite Element Atomic(%) 150 18.74 50.95 31.72 Totals 00.00 Fig. 5. The magnification of Zone B in Fig. 4
360 J. Zhong et al. / Journal of Materials Processing Technology 190 (2007) 358–362 Fig. 3. The magnification of Zone C in Fig. 4. Fig. 4. Fracture morphology of the C/C/SiC composite. thermal stress wound generate, which leaded to matrix cracking of composite. The micro-cracks could cause the deflection of the main crack when the material was subjected to load, which could improve the toughness of the composite. Fig. 5. The magnification of Zone B in Fig. 4. Fig. 6. The magnification of Zone B in Fig. 4. Fig. 7. The magnification of Zone A in Fig. 4. 3.2. Analysis of phase structure and chemical composition The XRD pattern showed in Fig. 8 indicated that the matrix completely consisted of the amorphous substances, which was similar to the results in Refs. [8,9]. Because the pyrolysis temperature was low and atoms did not have reactivity to crystallize, which resulted in the non-crystalline structure of the matrix. The EDS pattern which was obtained from the position marked in Fig. 9 and the analysis result were given in Table 2, respectively. It showed that the chemical compositions of matrix were mainly Si, C, N and C seemed excessive according to the stoichiometric Table 2 Elemental compositions of matrix of the as-prepared composite Element Weight (%) Atomic (%) C 34.03 49.54 N 15.02 18.74 Si 50.95 31.72 Totals 100.00 100.00
J Zhong et al. Journal of Materials Processing Technology 190(2007)358-362 36 room temperature 1300c wliw why strain, x 10 Fig 8. XRD pattern of matrix of the as-prepared composite Fig. 10. Stress vs strain curve for the as-prepared C/C/SiC composite. proportion of SiC. From this fact together with XRD pattern, it the as-prepared C/C/SiC composite. The stress-strain curve was suggested that the matrix formed by PlP did not consist of both at room temperature and 1300 C appeared ascending a mixture of carbon, SiC and Si3N4. It seemed to be a real com- zigzag, and that at room temperature exhibited more obvi- osite with Si-,C, and N-atoms as constituents. Consequently, ous non-linear behavior. As we known, the properties of the matrix was only considered as a Si-C-N composite, which material were determined by the structure. Zigzag character was C/C/SiC of the stress-strain curve probably is related to the lami nated structure of as-prepared composite showed in Fig. 4 3.3. Three-point-bending test When the crack propagated to the zone B, because of crack deflection, fiber debonding and pull-out, the crack propaga Three specimens were tested for room temperature and tion speed slowed down, and the stress increased with the 1300C, respectively. The test results dispersion was about increase of the strain until the crack crossed over the zone ten percent. Fig. 10 showed the typical failure behavior of B. With the increase of the stress, the damage in the co posite gradually increased. Damage implied the deformation in the macroscopic scale. When the damage accumulated to a certain extent, a pseudo-plasticity behavior appeared, in he way of zigzag stress-strain curves. It was obvious that the strain was very small both at room temperature and The stress-strain curve at 1300C displayed a typical fail ure process. The initial linear region AB corresponded to the elastic linear behavior of the material prior to any significant micro-failures. There was no damage generating in the com- posite at this stage. At an applied stress of point B there was a marked departure from linear behavior corresponding to trans- verse matrix cracking and to fiber/matrix debonding Region BC was believed to correspond to transverse crack multiplication 100um up to a saturation density. At an applied stress of point C fiber Electron Image 1 fractures and debonding which started before transverse matrix Spectrum 1 crack saturation continue until final failure [10]. The failure pro- cess at room temperature was similar to that of 1300C. But the flexural strength at room temperature was 121M was lower than 150 MPa at 1300C. The reason was that the residual thermal stress in the composite was partially relaxed at 1300°C. During the composite damaging process, there were several energy absorbed mechanisms which included matrix cracking, interface debonding, fiber pull-out and fiber fracture, these phe nomena did not occur in the monolithic ceramics. Owing to Full Scale 7411 cts Cursor. 0.000 damage process, the composite achieved more toughness com- Fig- 9. EDS pattern of matrix of the as-prepared composite paring with the matrix materi
J. Zhong et al. / Journal of Materials Processing Technology 190 (2007) 358–362 361 Fig. 8. XRD pattern of matrix of the as-prepared composite. proportion of SiC. From this fact together with XRD pattern, it was suggested that the matrix formed by PIP did not consist of a mixture of carbon, SiC and Si3N4. It seemed to be a real composite with Si–, C–, and N-atoms as constituents. Consequently, the matrix was only considered as a Si–C–N composite, which was C/C/SiC. 3.3. Three-point-bending test Three specimens were tested for room temperature and 1300 ◦C, respectively. The test results dispersion was about ten percent. Fig. 10 showed the typical failure behavior of Fig. 9. EDS pattern of matrix of the as-prepared composite. Fig. 10. Stress vs. strain curve for the as-prepared C/C/SiC composite. the as-prepared C/C/SiC composite. The stress–strain curve both at room temperature and 1300 ◦C appeared ascending zigzag, and that at room temperature exhibited more obvious non-linear behavior. As we known, the properties of material were determined by the structure. Zigzag character of the stress–strain curve probably is related to the laminated structure of as-prepared composite showed in Fig. 4. When the crack propagated to the zone B, because of crack deflection, fiber debonding and pull-out, the crack propagation speed slowed down, and the stress increased with the increase of the strain until the crack crossed over the zone B. With the increase of the stress, the damage in the composite gradually increased. Damage implied the deformation in the macroscopic scale. When the damage accumulated to a certain extent, a pseudo-plasticity behavior appeared, in the way of zigzag stress–strain curves. It was obvious that the strain was very small both at room temperature and 1300 ◦C. The stress–strain curve at 1300 ◦C displayed a typical failure process. The initial linear region AB corresponded to the elastic linear behavior of the material prior to any significant micro-failures. There was no damage generating in the composite at this stage. At an applied stress of point B there was a marked departure from linear behavior corresponding to transverse matrix cracking and to fiber/matrix debonding. Region BC was believed to correspond to transverse crack multiplication up to a saturation density. At an applied stress of point C fiber fractures and debonding which started before transverse matrix crack saturation continue until final failure [10]. The failure process at room temperature was similar to that of 1300 ◦C. But the flexural strength at room temperature was 121 MPa, which was lower than 150 MPa at 1300 ◦C. The reason was that the residual thermal stress in the composite was partially relaxed at 1300 ◦C. During the composite damaging process, there were several energy absorbed mechanisms which included matrix cracking, interface debonding, fiber pull-out and fiber fracture, these phenomena did not occur in the monolithic ceramics. Owing to damage process, the composite achieved more toughness comparing with the matrix material
J. Zhong et al. /Journal of Materials Processing Technology 190(2007)358-362 4. Conclusions References 1. The pyrolyzed substance of HMDS contained excessive car- II LCasas, MRElizalde, JM.MarTinez-Esnaola, Interface characterisation nd a small of and correlation with the creep behaviour of a 2. 5D SiC/C/SiC composite, Composites: Part A 33(2002)1449-1452. The phase composition was difficult to distinguish, primarily [2] X. Yongdong, C. Laifei, Z Litong, Y Hongfeng, Y Xiaowei, Mechanical considered as the Si-C-n compos properties of 3D fiber reinforced C/SiC composites, Mater. Sci. Eng. A 300 2. The structure of C/C/SiC composite was dense and homoge (2001)196-202 neous, but there were some pores between fibers and bundles, [3] V.I. Academician. Trefilov, Ceramic and carbon-matrix composites, Chap- man Hall London. 1995 whose sizes were less than 5 um and less than 100 pm, [4]P. Delhaes, Chemical vapor deposition and infiltration processes of carbon respectively. The thickness of the pyrocarbon interface was materials. Carbon 40(2002)641-6 non-homogeneous [5] M.F. Gonon, S. Hampshire, J.P. Disson, G. Fantozzi, A polysilazane pre- 3. The bending stress-strain curves were in the way of zigzag. cursor for Si-C-N-O matrix composites, J. Eur. Ceram Soc. 15(1995) 683-688 Fiber pull-out was observed on the fracture surface Flexural (6)M Qing-Song, C. Zhao-Hui, Z. Weng-Wei, H. Hai-Feng. Processing and strength of C/C/SiC composite at 1300C was higher than characterization of three-dimensional carbon fiber reinforced Si-O-Co that at room temperature. The main reason was that the resid- posites via precursor pyrolysis, Mater. Sci. Eng. A 352(2003)212-216 ual thermal stress in the composite was partially relaxed at [71 S. Zhou, S. Qiao, Study of flexural properties of C/C composite at high 1300°C. temperature, JMater Eng. 6(2001)16-18 [8]H Drost, M. Friedrich, R Mohr, E Gey, Nanoscaled Si-C-N-composite powders with different structures by shock-wave pyrolysis of organic pre- Acknowledgements cursors, Nucl. Instrum. Methods Phy Res B 122(1997)598-601 [9] L. Kleps, F. Caccavale, G. Brusatin, A. Angelescu, L. Armelao, LPCVD This research has been supported by the Foundation of silicon carbide and silicon carbonitride films using liquid single precursors, National Key Laboratory for Precision Hot Processing of Metals (10) 1M. Daniel. G.Anastassopoulos, The behavior of ceramic matrix fiber and Program for Changjiang Scholars and Innovative Research composites under longitudinal loading, Compos. Sci. Technol. 46(1993) 105-113
362 J. Zhong et al. / Journal of Materials Processing Technology 190 (2007) 358–362 4. Conclusions 1. The pyrolyzed substance of HMDS contained excessive carbon, followed by silicon, and a small amount of nitrogen. The phase composition was difficult to distinguish, primarily considered as the Si–C–N composite. 2. The structure of C/C/SiC composite was dense and homogeneous, but there were some pores between fibers and bundles, whose sizes were less than 5 m and less than 100 m, respectively. The thickness of the pyrocarbon interface was non-homogeneous. 3. The bending stress–strain curves were in the way of zigzag. Fiber pull-out was observed on the fracture surface. Flexural strength of C/C/SiC composite at 1300 ◦C was higher than that at room temperature. The main reason was that the residual thermal stress in the composite was partially relaxed at 1300 ◦C. Acknowledgements This research has been supported by the Foundation of National Key Laboratory for Precision Hot Processing of Metals and Program for Changjiang Scholars and Innovative Research Team in University. References [1] L. Casas, M.R. Elizalde, J.M. MarTinez-Esnaola, Interface characterisation and correlation with the creep behaviour of a 2.5D SiC/C/SiC composite, Composites: Part A 33 (2002) 1449–1452. [2] X. Yongdong, C. Laifei, Z. Litong, Y. Hongfeng, Y Xiaowei, Mechanical properties of 3D fiber reinforced C/SiC composites, Mater. Sci. Eng. A 300 (2001) 196–202. [3] V.I. Academician, Trefilov, Ceramic and carbon-matrix composites, Chapman & Hall, London, 1995. [4] P. Delhaes, Chemical vapor deposition and infiltration processes of carbon materials, Carbon 40 (2002) 641–657. [5] M.F. Gonon, S. Hampshire, J.P. Disson, G. Fantozzi, A polysilazane precursor for Si–C–N–O matrix composites, J. Eur. Ceram. Soc. 15 (1995) 683–688. [6] M. Qing-Song, C. Zhao-Hui, Z. Weng-Wei, H. Hai-Feng, Processing and characterization of three-dimensional carbon fiber reinforced Si–O–C composites via precursor pyrolysis, Mater. Sci. Eng. A 352 (2003) 212–216. [7] S. Zhou, S. Qiao, Study of flexural properties of C/C composite at high temperature, J. Mater. Eng. 6 (2001) 16–18. [8] H. Drost, M. Friedrich, R. Mohr, E. Gey, Nanoscaled Si–C–N–composite powders with different structures by shock-wave pyrolysis of organic precursors, Nucl. Instrum. Methods Phy. Res. B 122 (1997) 598–601. [9] L. Kleps, F. Caccavale, G. Brusatin, A. Angelescu, L. Armelao, LPCVD silicon carbide and silicon carbonitride films using liquid single precursors, Vacuum 46 (1995) 979–981. [10] I.M. Daniel, G. Anastassopoulos, The behavior of ceramic matrix fiber, composites under longitudinal loading, Compos. Sci. Technol. 46 (1993) 105–113.